University of Groningen

Investigation of the Nanoscale Morphology in Industrially Relevant Clearcoats of WaterbornePolymer Colloids by Means of Variable-Angle-Grazing Incidence Small-Angle X-rayScatteringVagias, Apostolos; Chen, Qi; ten Brink, Gert H.; Hermida-Merino, Daniel; Scheerder, Jurgen;Portale, GiuseppePublished in:ACS Applied Polymer Materials

DOI:10.1021/acsapm.9b00601

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Citation for published version (APA):Vagias, A., Chen, Q., ten Brink, G. H., Hermida-Merino, D., Scheerder, J., & Portale, G. (2019).Investigation of the Nanoscale Morphology in Industrially Relevant Clearcoats of Waterborne PolymerColloids by Means of Variable-Angle-Grazing Incidence Small-Angle X-ray Scattering. ACS AppliedPolymer Materials, 1(9), 2482-2494. https://doi.org/10.1021/acsapm.9b00601

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Investigation of the Nanoscale Morphology in Industrially RelevantClearcoats of Waterborne Polymer Colloids by Means of Variable-Angle Grazing Incidence Small-Angle X‑ray ScatteringApostolos Vagias,†,‡ Qi Chen,∥ Gert H. ten Brink,§ Daniel Hermida-Merino,⊥ Jurgen Scheerder,∥

and Giuseppe Portale*,†,‡

†Dutch Polymer Institute (DPI), P.O. Box 902, 5600 AX Eindhoven, The Netherlands‡Macromolecular Chemistry and New Polymeric Materials Group, Zernike Institute for Advanced Materials, and §Zernike Institutefor Advanced Materials, University of Groningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands∥DSM Coating Resins B.V. Sluisweg 12, Waalwijk 5145 PE, The Netherlands⊥DUBBLE Beamline at the ESRF, Netherlands Organization for Scientific Research, 71 Avenue des Martyrs, CS40220, 38043Grenoble, France

*S Supporting Information

ABSTRACT: Soft polymer colloidal water suspensions are extremely importantformulations for industrial applications such as water-based environmental-friendly coatings, paints, and adhesives. Homogeneity of the final coating at themicrometer and nanoscale is a crucial factor for optimal coating performance, suchas barrier properties against solvent permeation. Here, we investigated theremnant nanostructure in slot-die-coated micrometer-sized thick clear coatingfilms (clearcoats) of three different waterborne polymer colloids (pure soft, purehard, and soft/hard multiphase), commonly utilized as primers in paintformulations [Mader et al. Prog. Org. Coat. 2011, 71, 123−135], using variable-angle grazing incidence small-angle X-ray scattering (GISAXS) complementedwith cross-sectional atomic force microscopy (cs-AFM). After complete macroscopic drying, the coating films exhibit thepresence of residual nanostructure with characteristic distance (d*) smaller than the original particle size and even smaller(≪d*) heterogeneity dimensions. These nanostructural heterogeneities (i) develop due to partial particle coalescence, (ii) arepreferentially located close to the air−film interface and (iii) demonstrate the tendency to align perpendicular to the air−filminterface, implying vertical gradient in hydroplasticization effects having occurred earlier during film formation. The extent andsize of the nanostructural heterogeneities, driven by the slot-die coating application, strongly depend on the polymer chemistry(glass transition temperature, Tg) and the colloidal architecture. Last, solvent exposure has a significant impact on thenanostructure, causing the removal of these heterogeneities and leading to a more strongly coalesced film.

KEYWORDS: waterborne polymer coatings, acrylics, slot-die coating, variable-angle GISAXS, nanostructure, annealing, AFM,glass transition

■ INTRODUCTION

Waterborne polymer colloids are increasingly utilized as maincomponents in resin formulations for paint applications, barriercoatings, and anticorrosion products due to the reducedenvironmental and health hazards compared to solvent-basedresins.2−6 However, the distribution of colloids as a dispersedphase becomes challenging when using water as solvent.2,7 Anumber of microscale defects can develop during film drying.3,8

The interplay between the glass transition temperature (Tg) ofthe colloids and the drying temperature, as well as the sizedistribution of polymer colloids, can have a crucial impact onthe film’s structure and mechanical properties including thelateral and vertical segregation of nanostructural featureswithin the coating.7,9,10 Elucidating how variabilities inwaterborne polymer colloids, such as Tg, molecular weight,chemical composition, and colloidal architecture, can influence

the film structure is key in optimizing the final properties ofwaterborne coatings.11 Ideally, an optimal clearcoat film forpaint and protective coating applications (e.g., clearcoats/topcoats) should be free of defects, minimize the impact ofplasticizing agents such as water and ethanol, exhibit stronganticorrosive and solvent barrier properties,12 and demonstratesignificant resistance against “weathering effects”13 whenexposed to outdoor environment. To achieve these features,a clearcoat needs to balance between facile deformability/spreadability (low Tg) and high mechanical strength (highTg).

11,14 It is thus mandatory to noninvasively inspect thepossible presence of nanostructural heterogeneities in micro-

Received: June 30, 2019Accepted: August 23, 2019Published: August 23, 2019

Article

pubs.acs.org/acsapmCite This: ACS Appl. Polym. Mater. 2019, 1, 2482−2494

© 2019 American Chemical Society 2482 DOI: 10.1021/acsapm.9b00601ACS Appl. Polym. Mater. 2019, 1, 2482−2494

This is an open access article published under a Creative Commons Non-Commercial NoDerivative Works (CC-BY-NC-ND) Attribution License, which permits copying andredistribution of the article, and creation of adaptations, all for non-commercial purposes.

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meter-sized thick coatings and to investigate the influence ofparticle composition and architecture on those heterogeneities.The currently established drying mechanism of waterborne

latex suspensions involves the following stages:2,7 closepacking, deformation, and chain interdiffusion. The filmformation involves the competitive interaction between rateof water loss (convection), diffusion of polymer colloids,interfacial tension, capillary forces, and latex suspension’s (low-shear) viscosity. Imperfections arising during the packing anddeformation stages could affect the coating structure.15,16 Twotypes of constituents have been frequently assigned toheterogeneities: water17−19 and surfactant stabilizers.9,16,20−23

While information about the presence of water “pockets” incoatings from soft (Tfilm formation > Tg) polymer colloids withinthe submicrometer range has been reported,24 quantificationabout the exact characteristic length scales (distance and thesize) of such heterogeneities is less available. To study theexistence and the distribution of submicrometric heterogene-ities, one should aim at characterization techniques with broadspatial resolution, high surface sensitivity, and the possibility toprobe buried features. With spatial resolution of a fewnanometers, AFM can probe only the upper coating surface,while electron microscopy-based methods can be destructiveand under circumstances artifacts could be inevitable. More-over, the aforementioned local methods do not allow forsufficient statistics over the whole sample. Small-angle X-ray(SAXS) and small-angle neutron scattering (SANS) techniquesare optimal for characterizing particle morphology in polymercolloidal suspensions.25−27 These techniques have been alsoused to probe nanostructural heterogeneities in free-standingfilms either of hard, nondeformable polymer colloids28−31 orfor soft, deformable polymer colloids.32−34 Slot-die coating isthe key deposition method for application-relevant coatings(e.g., paint brushing or varnish application treatments). Itinvolves the deposition of a liquid formulation onto a substratetoward a controllable and uniform final film thickness.35 It isthus highly valuable to inspect the structure of coating filmsobtained by using slot-die coating, as a function of colloidsoftness, directly in their supported form without any furthermodifications. These technical challenges can be bypassedusing grazing incidence small-angle X-ray scattering (GISAXS),where the use of a 2D detector allows probing structuralvariations with spatial resolution between 1 and 1000 nmsimultaneously in the film plane and out-of-plane.36,37 WhileGISAXS is a well-established and powerful technique to studynanostructure in submicrometrically thin films,38−41 to the bestof our knowledge it has rarely been utilized to study muchthicker films (thickness of several micrometers), with theexception of one case on diblock copolymer films.42

Thicknesses between 1 and well above 10 μm are commonlyutilized for real applications such as waterborne polymerclearcoats (e.g., top-coating primers).43

In this study, we employ GISAXS to monitor the in-planeand out-of-plane order in the film nanostructure of industriallyrelevant waterborne coatings. These coatings are made of latexpolyacrylic nanoparticles with an approximate particle size of100 nm. The waterborne coatings in this work are used asclearcoats (clear finishes) for wood furniture, doors, andjoinery (window and door frames) as well as paint primers,43

their main function being to protect the substrates fromexposure to liquids such as water, coffee, and red wine (e.g.,alcohol). Our goal is threefold: (1) to inspect and quantify thedistribution of nanostructured heterogeneities within the film,

(2) to assess the influence of particle softness and architectureon those heterogeneities, and (3) to study the correlationbetween coating nanostructure and coating barrier properties.Synergizing structural information obtained from GISAXS withcomplementary cross-sectional AFM results, our workproposes a robust protocol to inspect the distribution ofnanometer-sized spatial heterogeneities across the film andperpendicular to its surface.

■ MATERIALS AND METHODSSynthetic Procedure. Three different polymer colloidal particles

were prepared by semicontinuous emulsion polymerization andnamed “S”, “H”, and “HS”, depending on their glass transition (seeTable 2). The following chemicals were used as supplied. n-Butylmethacrylate (n-BMA), n-butyl acrylate (n-BA), methyl methacrylate(MMA), and acrylic acid (AA) (Dow Chemical Company) were usedas monomers. Ammonium persulfate (United Initiators GmbH) wasused as initiator. Sodium bicarbonate and ammonia (Innophos andBrenntag Nederland B.V.) were used to regulate the pH. RhodafacRS/710E-30 (Solvay) was used as emulsifier. Proxel Ultra10 (ArchUK Biocides Ltd.) is added as preservative. The synthesis of thesingle-phase S and H particles is comprised of three steps, while oneextra step is required for the HS particles.

Step 1. A 2000 cm3flask equipped with a thermometer, a N2 inlet,

and an overhead stirrer was charged with demineralized water (685.9g), sodium bicarbonate (0.4 g), ammonia (0.8 g, 25 wt % solution),and Rhodafac RS/710E-30 (27.6 g). A first emulsified monomer feed,specific for each formulation, was prepared in feeding funnelsaccording to Table 1.

Step 2. For each formulation, a glass reactor was charged with asolution of ammonium persulfate (2.1 g), sodium bicarbonate (0.1 g),and Rhodafac RS/710E-30 (4.8 g) in demineralized water (42.9 g)and was heated to T = 85 °C. Once 85 °C was reached, 5 wt % of thefirst monomer feed was added, followed by the addition of anammonium persulfate solution (0.4 g for “HS”, 0.4 g for “H”, and 2.1g “HS” in 4.7 g of demineralized water). The reaction was firstallowed to reach the peak temperature (T ∼ 89 °C), and then the restof the first monomer feed was added in 120 min for “HS” and in 180min for both “S” and “H”. After the feeding period, the reactionmixture was left to react further at T = 85 °C for another 45 min.

Step 3 (Only for HS). A second emulsified monomer feed wasprepared by mixing n-butyl methacrylate (n-BMA, 50.8 g), methylmethacrylate (MMA, 149.2 g), acrylic acid (AA, 10.5 g), sodiumbicarbonate (0.3 g), Rhodafac RS/710E-30 (7.2 g), and demineralizedwater (94.3 g) until a stable monomer feed was obtained. After the 45min holding period, this second monomer feed was added in 60 mintogether with the remainder of the ammonium persulfate solution,and the reaction mixture was left to react at T = 85 °C for another 30min.

Step 4. The reaction mixture was allowed to cool to T = 23 °C, thepH was adjusted by adding a dilute ammonia solution (≈4 wt %), andProxel Ultra10 was added (5.0 g). The solid content was measured byusing a Mettler Toledo HB-43S moisture analyzer using 105 °C asdrying temperature and was adjusted to 39 wt % by addition of

Table 1. Composition of the Examined PA-Based Materials(All Values in grams)

S H HS

n-butyl methacrylate (n-BMA) 428.7 169.3 300.0n-butyl acrylate (n-BA) 200.5 140.3methyl methacrylate (MMA) 37.5 497.3 26.3acrylic acid (AA) 35.1 35.1 24.6sodium bicarbonate 0.5 0.5 0.2Rhodafac RS/710E-30 15.3 15.3 8.1demineralized water 236.0 236.0 165.2

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demineralized water. Finally, the batch was filtered over a 75 μm clothand collected. The conversion was measured by GC and was >99.9%.Formulation Characterization. The sample glass transition (Tg)

was measured by differential scanning calorimetry (DSC). Thermaltransitions of the synthesized formulations are reported in Figure S1,and associated parameters (Tg and Cp) per formulation aresummarized in both Table 2 and Table S1. The relevant physical

properties for each formulation are summarized in Table 2. Theparticle size was measured by dynamic light scattering (DLS), and thesize distribution was evaluated by using the CONTIN algorithm. ATEM image for the HS particles is reported in Figure S2. The particlesize and architecture were also studied by solution small-angle X-rayscattering (SAXS) experiments (Figure 1 and Figure S3). SAXS was

performed at the MINA diffractometer of the University ofGroningen, equipped with a Cu rotating anode (X-ray wavelength λ= 1.5413 Å/energy of 8 keV) and using a Bruker Vantec 2000 2Ddetector placed 3 m away from the sample. An amount of 70 μL ofeach suspension was loaded into a 1.5 mm thick glass capillary. Thecapillary was flame-sealed to prevent water evaporation, and theexposure time was of the order of 10−30 min depending on thepolymer concentration. SAXS images were normalized by theexposure time and were integrated by using the Fit2D software.The scattering angles 2θ were calibrated using the position of knowndiffraction peaks from a standard silver behenate powder, and theSAXS intensity was finally reported as a function of the modulus ofthe scattering vector q = (4π/λ) sin θ.Film Preparation. Films were prepared at T = 23 °C by slot-die

coating onto sodalime glass slides (75 mm × 25 mm dimensions,

Menzel). An amount of 160 μL from aqueous HS, S, and Hsuspensions was coated on the glass by using aluminum slot dies withdifferent gap clearances (10 and 120 μm), and different startingpolymer concentrations in water were used to achieve differentuniform film thicknesses (Figure S4). Water evaporation was beenstudied by recording the weight loss of the different HS, S, and Hsamples over time (Figure S5).

Film Characterization. Atomic Force Microscopy (AFM)Imaging. AFM measurements were performed in at least three dryfilms per colloidal formulation and at three different locations per filmseparated by several micrometers one from each other. AFMtopography images were acquired at the film−air interface, with aNanoScope V multimode atomic force microscope (Bruker NanoSurfaces, Santa Barbara, CA) using silicon cantilevers with resonancefrequencies of 300−400 kHz (model: TESP, Bruker Nano surfaces).

AFM Imaging of Coating Cross Section (cs-AFM). The crosssection of the HS coating, applied on thin (150 μm) borosilicate glassslide, was prepared by using a Leica EM TIC3X ion beam millingsystem by exposing the cross section to argon ions with accelerationvoltage of 4 kV and gun current of 1.6 mA for 8 h under ambienttemperature. We prepared one single HS sample using a triple-ionbeam miller. The milled area was ∼5 mm × 0.5 mm (the totalthickness of the glass slide + coating). The ion beam milled samplewas mounted onto an AFM vertical sample holder. The cs-AFMimages were obtained under ambient conditions in tapping mode witha NanoScope V multimode atomic force microscope (Bruker NanoSurfaces, Santa Barbara, CA) using silicon cantilevers with resonancefrequencies of 300−400 kHz (model: TESP, Bruker Nano surfaces).In addition, a single noncolloidal-based atactic polystyrene film wasprepared and characterized by cs-AFM. PS with 370K molecularweight (Polymer Source Inc.) was spin-coated under 400 rpm for 60 sfrom a 6 wt % toluene solution and annealed at T (= 140 °C) > Tg,PS(∼100 °C)44 for 1 h to repair defects stemming from spin-coating.The roughness of the analyzed cross sections was determined from 10different randomly selected areas (2 μm × 2 μm) at the same depth.In total, micrographs at eight different depths from the air−filminterface have been acquired.

Grazing Incidence X-ray Scattering (GISAXS). GISAXS experi-ments were conducted at T = 23 °C, between 2 and 4 days followingfilm preparation, at the Dutch-Belgian beamline (DUBBLE, BM26B)at the ESRF, Grenoble, using 12 keV irradiation energy (λ = 1.033Å).45,46 In total, at least three different films per colloidal formulation(H, S, and HS) have been examined. The X-ray beam was focused atthe sample position and had dimensions of about 300 μm(perpendicular to the film surface) × 1000 μm (along the filmsurface). To resolve scattering features of the relevant morphologiesin the examined systems, we used a sample-to-detector distance (S-to-D) of 7 m. The GISAXS 2D patterns were recorded on a noiseless,solid-state Pilatus 1M Dectris detector, with 981 (laterally) × 1043(vertically) pixels and pixel size of 172 μm × 172 μm as a function ofqy and qz (the component qx along the beam direction can beneglected here). The moduli of the in-plane and out-of-planescattering wavevectors are respectively

πλ

α ψ= ikjjj

y{zzzq

2cos( ) sin( )y f (1)

πλ

α α= [ + ]ikjjj

y{zzzq

2sin( ) sin( )z f i (2)

where ψ is the in-plane scattering angle in the direction parallel to thefilm surface, αi is the incident angle of the X-ray beam, and αf is theexit scattering angle in the vertical direction perpendicular to the filmsurface. The full measured in-plane range was qy = 0.02−1.7 nm−1

(i.e., 4−290 nm), and the out-of-plane one was up to qz = 1.3 nm−1.Note that the GISAXS patterns presented in this article are a zoomed-in view of the larger acquired patterns. All the acquired GISAXSpatterns have been normalized for the incoming beam intensity. Thenominal critical angles αc of the polymer films and of the sodalimeglass substrate are ∼0.1°−0.11° and 0.15°, respectively. The samplesurface was aligned with respect to the beam direction by using a

Table 2. Main Physical Properties of the Examined PA-Based Materials

property HS (hard−soft) S (soft) H (hard)

solid content (%) 39.0 39.1 39.1pH 7.0 6.8 8.1particle radius by DLS, R (nm) 53 ± 2 54 ± 1 50 ± 2Tg, midpoint (°C) 4.6 and 96.1 5.0 94.9Cp [J g

−1 °C−1] 0.20 and 0.08 0.29 0.32

Figure 1. SAXS intensity curves normalized with respect to theacquisition time, I(q) (counts s−1), vs wavevector q (nm−1) foraqueous suspensions of the different formulations at 39 wt % solidcontent: HS (black), S (green), and H (blue). The vertical dashed-dotted arrows point to the first minima position of the particle formfactor, and the short solid arrows denote the peak position q* fromthe structure factor S(q) for each color-matching curve.

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high-resolution HUBER circular segment goniometer. The estimatederror on αi was ±0.05° (as determined from the uncertainty incalculating the center of the transmitted beam profile scans from thephotodiode embedded in the beamstop and the position of thereflected X-ray beam on the detector). The incident angle αi wasvaried from close to (αi ∼ αc) to well above (αi > αc) the critical angleαc of the polymer to probe the sample nanostructure at differentdistances starting from close to the air−film interface down to theglass substrate. The total accumulation time for each αi was 300 s.Because of the particular nature of the examined films (waviness at thefilm−air interface and relatively thick, micrometer-sized films), thetransmitted scattering signal increases over the reflected one, and theYoneda47 peak appears broadened along qz. Because of thisbroadening, the I(qy) vs qy intensity cuts have been better computedat qz positions slightly higher than the position of the Yoneda peakheight (e.g., up to 20−30 pixels above the Yoneda peak position) tomore clearly observe the scattering peaks. The horizontal cuts havebeen obtained by averaging the intensity of ±5 adjacent rows alongthe qz-axis (Δαf = 0.012°).

■ RESULTS AND DISCUSSIONThe water-based colloidal suspensions of each formulationhave been first characterized by transmission SAXS (Figure 1).The SAXS curves for the three different colloidal suspensionsshow several oscillations at high q values typical for the particleform factor together with a peak in the region q < 0.1 nm−1.48

The position of the scattering minima matches the oneexpected for spherical particles.The minima are damped due to polydispersity in particle

size, denoted here as σrel = σ/R, with σ being the standarddeviation of the particle size distribution function (Schulz−Zimm in this case). Analysis of the first minimum position andfitting of the SAXS curves (Figure S3 and Table S2) providesthe following spherical particle radii Rp,S = 50 nm (σrel = 0.07),Rp,H = 45 nm (σrel = 0.15), and Rp,HS = 46 nm (σrel = 0.16),slightly smaller than the DLS results reported in Table 2. Sucha discrepancy is often reported between these two methods.49

The TEM micrograph for HS particles (Figure S2) confirmsthe size polydispersity observed by SAXS. Interestingly, whilethe SAXS curves of the pure H and S suspensions can besuccessfully described by using a simple homogeneous spheremodel, the SAXS curve for the multiphase HS required the useof a two-phase model (see Figure S3 and “SAXS Modeling ofSuspension Data” section in the Supporting Information). TheHS particle form factor can be well described by a concentriccore−shell spherical model (details provided in the SupportingInformation). The remaining minor discrepancies between themodeled curve and the experimental data could be attributedto several reasons such as deviation from a spherical shape,incorrect assumption of the size distribution function,inhomogeneities in the polymer shell, or the presence of asurfactant corona, not considered here. The scattering peakslocated at q* is generated by the maximum of the structurefactor describing the spatial correlation among particles and isrelated to the average interparticle distance (dinterparticle = 2π/q*).48 Variability of the interparticle distance in suspensions(dinterparticle,H = 94 nm < dinterparticle,S = 115 nm < dinterparticle,HS =125 nm) may probably reflect differences in the extent ofinterparticle interactions stemming from pH variations (Table2) or differences in the relative amounts of (charged)surfactant stabilizer and ionization degrees of different weakorganic polyacids.49,50

By use of slot-die coating, the deposition of high-qualitycoating films with controlled thickness and high homogeneityalong the in-plane direction has been mastered (Figure S4).

SEM images show the achieved uniformity in both filmthickness and topography (see Figure S5). The thicknessclosely relates to the thickness of single layer of paintbrushcoatings.51,52 A detailed stylus profilometry study suggested aroot-mean-squared waviness Wq = 1.5 μm for such thick filmsand showed that thickness nonuniformities at the film edgesare present, but their contribution is very limited as theyrepresent not more than 15% of the total film width distributedon the two edges (Figure S5). The availability of such goodquality films is fundamental for the GISAXS study presentedhere.AFM images acquired on the surface of ∼30 μm thick films

are reported in Figure 2 and suggest partially coalesced

particles in the case of HS film with a root-mean-squareroughness Rq,HS = 2.6 ± 0.4 nm, about 4 times larger than forfilms of the soft S analogue, Rq,S = 0.6 ± 0.3 nm.AFM has an excellent spatial resolution but is limited only to

topography information at the air−film interface. To probe thenanostructure close to the air−film interface as well as withinthe film and across the film thickness, we utilized variable-angleGISAXS analysis. GISAXS is a powerful technique to study avariety of structures found in supported thin films.37,53 Itsapplication to thick (>1 μm) films is more challenging but stillpossible. Tuning the incidence angle (αi) of the X-rays withrespect to the sample surface allows one to change the X-raypenetration depth (ξp) inside the film.42 This approach wasefficiently used in the past to obtain depth-resolved structuralinformation by using GISAXS on block copolymer thick filmsand carbon-based/nanoparticle films.54−56 For a given samplewith X-ray refractive index n = 1 − δ − iβ, the penetrationdepth ξp (depth at which the beam intensity is attenuated by1/e) depends on the material’s critical angle (α δ= 2c ), theX-ray wavelength (energy of the X-ray photons), and theincident angle αi (eq S7). Typical ξp(αi) vs αi curves computedfor polyacrylic-based films that closely represent thecomposition of our coatings at an X-ray energy of E = 12keV are plotted in Figure S6. The calculated values for ξp(αi)show that in principle, depending on αi, the ξp can be tunedfrom about 5−10 nm up to tens of micrometers. However, itshould be considered that the presence of surface defects,surface undulations, and menisci limits the accuracy in thedetermination of αi and hinders the precise calculation of ξp,especially at low αi. This is particularly true for thick polymerfilms as the ones used here. Thus, the actual minimumpenetration depth at αi ≈ αc is expected to be well above 10nm, and the ξp values for αi > αc mentioned here should beconsidered as nominal values for the intended depthresolution. Moreover, as hard X-rays (E > 6 keV) were used,ξp increases quite rapidly for αi > αc. Thus, it is quite difficult

Figure 2. AFM height topography 1 μm × 1 μm images: HS (leftpanel) and S (right panel) films.

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to fine-tune ξp in the range from 20 nm to 1−2 μm. We havethus acquired GISAXS data at αi ≤ αc (αc,H = 0.11°; αc,HS = αc,S

= 0.105°) and well above it (αi > 0.15°). Figure 3 shows acollection of the GISAXS patterns as a function of increasing αi

for the three investigated coatings with similar thickness (h ∼30 μm).The characteristic features present in such GISAXS patterns

are explained in detail in Figure S7. First, the GISAXS results

for the multiphase HS film (Figure 3, top row) are discussed. Astrong change on both the pattern appearance and thescattering intensity is observed with increasing αi. For αi ≈αc = 0.1°, the GISAXS pattern shows two clear symmetricanisotropic scattering signals with respect to qy = 0 nm−1. Asexpected, the intensity of these signals is enhanced close to theYoneda peak position along the qz direction.

47,57 The signalintensity decreases significantly with increasing αi. Interest-

Figure 3. GISAXS patterns as a function of incident angle αi for the different films slot die coated on sodalime glass substrates: HS (top row); S(middle row); H (bottom row). The intensity scales as follows: from 0 to 7 × 10−4 (HS), 0 to 5 × 10−4 (S) and then for (H), from 0 to 2 × 10−1

(αi = 0.08°; αi = 0.17°), 0 to 5 × 10−2 (αi = 0.22°), 0 to 10−2 (αi = 0.43°) and 0 to 2 × 10−3 (αi = 0.59°). At the lowest row (αi = 0.59°, Hcoatings), the specular signal is shown to the side as it was present at higher qz than the one displayed for all the other scattering patterns.Concerning HS and S coatings, the nominal penetration depth values (ξp(αi)) associated with the used incident angles (αi) are ∼10 nm (αi = 0.1°);∼10 μm (αi = 0.15°); ∼13 μm (αi = 0.17°); ∼16 μm (αi = 0.2°); ∼19 μm (αi = 0.22°); ∼35 μm (αi = 0.37°); ∼47 μm (αi = 0.5°); ∼50 μm (αi =0.53°). Similarly, for H coatings: 6 nm (αi = 0.08°); 11 μm (αi = 0.17°); 16 μm (αi = 0.22°); ∼35 μm (αi = 0.43°); ∼45 μm (αi = 0.59°).

Figure 4. Intensity cuts I(qy) (counts s−1) vs qy for (a) the HS and (b) the S films as a function of αi. The cuts stem from the GISAXS scattering

patterns of the respective films on sodalime glass substrates. The solid arrow points to the direction of increasing (αi), and the dashed arrow pointsto the peak position q*(αi) of the characteristic correlation length, d*(αi). The gray-shaded rectangular region denotes the beamstop position.

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ingly, for αi > 0.25°, the scattering peaks are not visibleanymore. Moreover, edge effects on GISAXS scatteringpatterns (Figure S8, shown for HS) can be neglected. Althoughwith lower scattered intensity by ca. 6 times, the GISAXSpatterns for the soft S film (Figure 3, middle row) show thesame trend with increasing αi. In contrast to the multiphase HSand to the soft S films, the GISAXS patterns from the nonfilmforming hard H colloidal formulation (Figure 3, lower row)exhibit strong scattering signal at all αi. As a consequence ofthe hard colloidal particle nature, the qualitative aspects of theGISAXS patterns for the H films resemble closely the ones forpolydisperse uncoalesced colloidal spheres.58,59 To quantita-tively analyze the GISAXS patterns, we performed horizontalcuts of the scattered intensity I(qy) along the in-plane qydirection at qz values slightly above the Yoneda peak heightposition (see Figure 4a for HS, Figure 4b for S, and Figure S9for H). For a complete analysis, we also report the I(qy) cuts atthe qz values of the specular beam position for coatings of allthree formulations in Figure S10, which show similar behavior.The presence of the symmetric scattering peaks for HS and

S is attributed here to the existence of a nanostructure whosein-plane characteristic correlation length/distance, d*(αi) =2π/q*y,max, is ∼80 nm, as estimated from the peak positionalong the qy direction (q*y,max = 0.078 nm−1 at αi = 0.10° forboth HS and S). d*(αi) seems to depend on the αi, and it wasfound to increase monotonically with increasing αi in the caseof the multiphase HS coating films, but not in the case of soft Sfilms (Figure 5). This distinct trend might imply synergy of

shearing effects and also different net interplay of (hydro-phobic, electrostatic, and hydrogen-bonding) interactionsbetween the polymer constituents and the soluble additivesdue to matrix topography differences, but also different particlesintering mechanism, pointing to a more regular packing forthe S sample (see the Drying-Deformation Calculations andSintering Mechanism section). The more regular packing ofthe S sample is also evidenced by the presence of a secondhigh-order scattering peak observed in Figure 4b. Conversely,the absence of higher order peaks in the I(qy) cuts reported inFigure 4a suggests a large extent of structural disorder for theHS films. Interestingly, the peak intensity is strongly enhancedand the peak shape is significantly different with respect to theanalogous SAXS peak in the aqueous suspensions (see Figure

1). This difference in peak characteristics (both position andshape) between solution SAXS and GISAXS together with thedisappearance of the intensity oscillations characteristic for thespherical particle shape suggests that the colloidal particleshave deformed and coalesced (to some extent) during drying.We thus attribute the origin of the GISAXS scattering peaks inboth HS and S films to the spatial correlation betweennanostructural heterogeneities generated by incompleteparticle coalescence and showing substantial electron densitydifference with respect to the polymer matrix. Remarkably,these nanostructural heterogeneities seem to exhibit certaindirectionality along the direction perpendicular to the filmsurface. The scattering signals for the HS and S film tend to fallon a ring but appear more focused and often elongated alongthe horizontal qy direction, suggesting a preferential orientationof the scattering entities along the vertical direction (Figure 3).These rings may suggest the presence of anisotropic structurestilted throughout the plane of the film, as it will be confirmedand discussed below on the basis of GISAXS simulations.Differently from the HS and S films, the intensity oscillationstypical of the particle form factor are still present for the hardnon-film-forming H film, and the peaks in the respective I(qy)cuts stem from a different reason (Figure S9): they representnanostructural spacing within a randomly packed array ofnoncoalesced and nondeformable hard colloids. Moreover, forthe H coating, the I(qy) does not present drastic changes incurve shape as a function of increasing αi compared to HS andS, apart from the shift in the peak structure factor (Figure S9).It should be noted that for very high αi the signal is mostlyconcentrated along the specular (higher qz) region, rather thanat the Yoneda (relatively smaller qz) refraction region, as aresult of the large sample roughness (Figure S9 as compared toFigure S10c). The variable-angle GISAXS peak intensity isproportional to the amount (number and size) and thescattering length density contrast of the nanostructuralheterogeneities in the coating film. According to ourinterpretation, larger peak intensity would imply moreincomplete particle coalescence and hence less extent ofchain interdiffusion between the polymer colloids, whiledisappearance of the scattering peak indicates a large (if notcomplete) degree of particle coalescence. As mentioned above,the intensity for the soft S film is ∼6 times lower than HS(Figure 3, top vs middle row). This can be understood by theexpected higher degree of particle coalescence for the softercolloidal S system (see also the “GISAXS simulations” sectionin the Supporting Information). Notably, despite the low Tg ofthe S sample (Table S1), the degree of coalescence is not 100%although being larger than the one for the multiphase HSformulation. The crossover from a structured upper part to anonstructured lower section of the film, indicated by thedisappearance of the scattering peak, was found at αi ∼ 0.20°−0.25°, corresponding to ξp ≈ 18 μm (Figure S6). A significantincrease of the average distance between the remainingheterogeneities d*(αi) with αi is observed for the HS film asa result of the larger degree of particle coalescence deeper inthe film (see Figure 5).To further verify that these heterogeneities are located inside

the film and not only at the air/polymer interface, we havecompared the GISAXS intensity cuts with the transmittedSAXS data for free-standing HS and S sample (see FigureS11a,b). The close agreement in both intensity and form factorcontribution between transmission SAXS experiments com-pared to the analogous GISAXS at αi = 0.15° (nominal ξp ∼ 10

Figure 5. Heterogeneity spacing d*(αi) as a function of incidentangle, αi, for 30 μm (solid black rhombi) thick films of HS and for 30μm (solid green triangles) S films. The lines are drawn to guide theeye. The error bar denotes the standard deviation associated with theuncertainty in determining the scattering peak positon (uncertainty inbeam center determination and detector resolution) per eachmeasurement run at a certain αi.

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μm) confirms that these nanostructures are indeed locatedinside the bulk subsurface region of the coating and extendseveral micrometers from the surface down to the bulk of thefilm. Furthermore, the surface structure probed by Fouriertransform of the AFM images via the power spectral densityplot (see Figure S11c) is quite different from the GISAXSresults, confirming that already at the lowest αi GISAXSpenetrates certain layers of particles.To estimate the average dimension of the nanostructural

heterogeneities, we attempted simulations of the GISAXSintensities. Simulation of GISAXS patterns from disorderedthick films (i.e., many layers of scattering objects) in theframework of the distorted wave Born approximation60 is aquite challenging task. We have thus opted here for simulationsand fits of the horizontal I(qy) intensity cuts only (see FigureS12 and Table S3 for a summary of the extracted structuralparameters) to qualitatively support the presented experimen-tal results. The horizontally scattered intensity profile I(qy) forthe H film is successfully described by an ensemble ofrandomly packed polydisperse spherical particles. This isexpected due to the hard nature (high Tg) of the respectivecolloidal H particles. Conversely, the I(qy) intensity profile ofsoft S sample looks different and can be well described using anensemble of narrow, vertically aligned, randomly packedcylindrical objects. The average cylinder diameter in the Sfilm represents the average cross section of the heterogeneitydomains Dheterog,S ∼ 8 nm. The calculated average distancebetween heterogeneities is about d*S,sim ∼ 72 nm, in agreementwith the value d*S,exp ∼ 80 nm obtained from the peakposition. These simulation results are in agreement with thelarger degree of coalescence of the soft S particles, leavingbehind small, well-separated elongated nanodomains with asignificant degree of vertical alignment. The intensity for themultiphase HS sample can be also described using a randomlypacked ensemble of cylinders (see Figure S12c−e), but with anaverage equivalent diameter of Dpolymer,HS ∼ 70 nm andcalculated average distance of d*HS,sim ∼ 73 nm, in agreementwith the experimentally determined average distance betweenheterogeneities of d*HS, exp ∼ 80 nm. The agreement betweenthe simulated curve and the experimental I(qy) intensity plotsfor the HS film is worse than for the H and the S films,suggesting that the morphology of the HS film is morecomplex. The simulations suggest that the scattering features inthe HS film are most likely composed by partially coalescedparticles with some degree of vertical alignment, leading toelongated polymeric structures. For the H and HS films, theheterogeneity domain size can be roughly estimated from theresidual space between the polymer domains (interstitialspace), using as input values the experimentally determinedaverage polymer domain distance between heterogeneities ofd*HS, similar to coordination of cations between anions61 (seethe “GISAXS simulations” section in the SupportingInformation). The simulation results suggest thus a larger“nanoporosity” of the HS film (and of course of the H film)with respect to the S film, in agreement with the experimentallyretrieved GISAXS scattering intensity. These nanostructuralheterogeneities can contain nonbound material as well as someresidual water. AFM tests conducted on as-prepared, water-exposed, and thoroughly rinsed S films highlight the initialpresence of surfactant on both the surface and inside the film,which can be driven to the coating surface upon waterexposure and eventually washed away (see Figure S13). Theresidual water content could be assessed by TGA analysis (see

Figure S14). Interestingly, the 1% water content measured byTGA for the S film is comparable with the heterogeneities’volume fraction estimated from the GISAXS simulations(ϕheterog ≈ 0.01; see Table S3). On the contrary, the residualwater content for the HS film is clearly lower than theestimated heterogeneity volume fraction (ϕheterog ≈ 0.23; seeTable S3). Thus, in percentage, the heterogeneities of the Sfilm are more filled by water than the HS heterogeneities, andthe latter (HS) are more likely to be open structures filled byair (see below for further discussion). Although qualitative,these simulation results clearly show that GISAXS can capturethe nanostructural difference between coatings of differentchemistry/morphology and provide a rough, first, estimate ofthe heterogeneities size and the coating nanoporosityotherwise difficult (if not impossible) to measure with othertechniques. Moreover, these heterogeneities at the nanoscalecould be utilized as a predictive caliber for the performancebehavior of waterborne latex coatings at the macroscopic scale.The vertical cuts I(αf) vs αf extracted from the 2D GISAXSimages at similar αi (∼0.15°) for the different samples havealso been analyzed and are reported in Figure S15. While boththe S and HS films show evidence for a clear Yoneda peak at anexit angle of αf ∼ 0.1°, the intensity cut of the H film onlyshows a broad scattering peak as a result of the large roughnessand the colloidal particle nature of this film. Interestingly, theHS sample shows a second Yoneda at αf ∼ 0.07°, suggestingthat this film may contain a region with large roughness orporosity at the polymer/air interface, in contrast to the Ssample. This observation is in line with the greater surfaceroughness of the HS film (see Figure 2). Relying on theexperimental GISAXS patterns, the GISAXS simulations, andthe thermogravimetry (Figure S14), we propose a schematicview of the film nanostructure (Figure 6).

According to our vision about the coatings’ nanostructure,we expect that the macroscopic barrier properties againstsolvent permeation4,12 should link to the coatings’ “nano-porosity” and is expected to increase as S < HS < H. To provethis conjecture, we conducted macroscopic staining tests onHS and S coatings.

Staining Experiments. Macroscopic staining tests havebeen conducted against wine and coffee permeation. As visible

Figure 6. Envisionedbased on the experimental GISAXSpatternsnanostructures from surface overview (upper line) andside view (lower line) for left to right: the soft/hard multiphase (HS),the soft (S, Tg ∼ 5 °C), and the hard (H, Tg ∼ 95 °C) films. Greenand gray colors denote the polymeric domains from the soft phase Sand the hard phase H, respectively. The blue color in the HS and Scoatings denotes the presence of residual water and other nonboundmaterials.

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from Figure 7, the impact of both coffee and wine staining ismuch stronger for the HS coating compared to the S coating,

suggesting a strong difference in coating barrier properties(details about these tests are reported in the Figure 7 caption).The stronger staining effects for HS coatings agrees well

with the suggested higher (open) porosity of the HS filmderived by the GISAXS analysis. Thus, the presence ofdispersed hard phase in the HS polymer colloid formulationfrom one side increases the mechanical strength and on theother hand seems to diminish their solvent-proof property12

and to strongly enhance out-of-plane heterogeneity forma-tion62 compared to the single-phase soft S film.

Cross-Sectional AFM (cs-AFM). To (1) probe that thedepth-resolved GISAXS trend is not technique-specific and to(2) provide additional evidence for partial coalescence, we alsoconducted cross-sectional AFM (cs-AFM) experiments onmicrometer-sized thick films of HS on a glass substrate (Figure8). We compare the root-mean-squared roughness (Rq)measurement as a function of distance from the air−polymerinterface obtained from cs-AFM with the trend in A(αi) vs αi

from GISAXS. Each individual Rq value is determined from a 2μm × 2 μm area (dashed 2 μm × 2 μm white box, Figure 8b,upper panel) related to a 50 μm × 50 μm AFM largermicrograph as shown below. The Rq was averaged at eightdifferent locations, at the same distance from the air−filminterface, along the coating cross section. The cs-AFM resultsdepict a monotonic decay of Rq from 8 to 3 nm with increasingdistance from the air−polymer interface. This finding is in linewith the experimental GISAXS evidence for structuralheterogeneities vanishing from the air−polymer interfacetoward the glass−polymer interface. To exclude the possiblepresence of artifacts induced by the ion milling process, weperformed a control cs-AFM experiment on a noncolloidal-based sample, namely atactic PS (Mw = 370K). The PS film(∼5 μm thick) that had been thermally annealed above its Tg

to repair for holes and defects exhibits negligible Rq, suggestingthat possible pitfalls, if any, during the ion milling preparationof cross-sectional AFM can be safely ruled out. Thus, theresults from both GISAXS and cs-AFM clearly suggestincomplete coalescence of the HS polymer colloids, withheterogeneities unevenly distributed across the film thickness.Hence, this structural anisotropy presents evidence in favor of

Figure 7. Macroscopic staining experiments on HS coating (left) andS coating (right): coffee stain effect (a, b); wine stain effect (c, d). Forthe cases of coffee and wine test, the coatings had been prepared usingthe aluminum slot die and after drying were kept in contact with asoaked cotton of the respective liquid for 3 h before they gotmacroscopically inspected. A 50 mm × 25 mm rectangular piece ofcotton was soaked in the respective liquid. Scale bar: 20 mm.

Figure 8. (a) Consecutive cross-sectional AFM micrographs (5 μm × 5 μm) from the air−film interface (top) down to the glass-film interface(bottom) for HS. Smaller (1 μm × 1 μm) boxes from the corresponding (color-matching) 5 μm × 5 μm images are also shown. (b) A 50 μm × 50μm cross-sectional AFM image, denoting from left to right, air, HS film, and glass. The 2 μm × 2 μm white squares indicate the areas where theRMS roughness (Rq) was calculated. Estimated values of the RMS roughness (Rq) as a function of the distance from the air−polymer interface arealso shown: HS (black squares) and PS (Mw = 370K) noncolloidal polymeric film prepared by spin-coating (blue circles). The error bar denotesthe standard deviation of the mean Rq based on eight different locations inside the coating but measured at the same distance from the air−filminterface.

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a vertical gradient in particle concentration in the HS filmcoatings, in agreement with simulations on latex coatings.23

Concerning the HS film, we speculate that the difference inRMS roughness between surface AFM (Rq,HS ∼ 2 nm) and cs-AFM few micrometers below the air−surface interface (Rq,HS ∼8 nm) could be linked to nonbound species from the emulsionpolymerization that accumulate at the top surface, similarly towhat observed for the S coatings (see Figure S13), althoughpreliminary Raman data did not prove clear chemical signatureof a different chemical species on top of the HS coatings.During evaporation, this (nonbound) water-soluble materialincluding surfactant will end up in the nanovoids and canexude to the coating surface through these nanoheterogene-ities.20,63

Solvent Exposure and Annealing. Annealing protocolsare frequently employed in films of polymer colloids to reducethe amount of heterogeneities and increase resistance againstsolvent permeation.31,64−66 To assess the impact of solventannealing on the nanostructural heterogeneities, we prepared apair of 30 μm samples (one HS coating; one S coating) whichwe exposed to ethanol (EtOH). Before recording GISAXSpatterns, the samples had been submerged in EtOH for 1 h andthen were fully dried for at least 1 h at ambient conditions. Asshown in Figure 9a−d (and Figure S16), EtOH treatment has

a drastic influence on the film nanostructure regardless of thepresence of high Tg domains in the HS film, and theheterogeneities are smeared out upon EtOH exposure, asevidenced by loss of GISAXS signal for both HS and S films.Weathering tests on protective coating performance include

the impact of aging.13 Interestingly, the heterogeneities are stillpresent in a 30 μm HS coating after aging at ambientconditions for 12 months. Τhose heterogeneities eventually getsmeared out upon thermal annealing T = 150 °C for 1 h,proven by loss of GISAXS signal for the HS film (see Figure9e,f and Figure S16c). It is clear that thermal annealingplasticizes the film and promotes expulsion of the embedded

cosolutes (water, ions, and surfactant molecules) that getentrapped in the heterogeneities.67

Drying-Deformation Calculations and SinteringMechanism. The differences in nanostructure revealed byGISAXS for the three studied coating formulations may belinked to differences in the sintering mechanism. The Pecletnumber of the latex particles is estimated to be PeNP = LwetE/DNP ∼ 4 by considering a diffusion coefficient for the latexnanoparticles of DNP ≈ 4.6 × 10−12 m2 s−1 (estimated using awater viscosity of μwater = 8.9 × 10−4 Pa·s at Τ = 23 °C and aparticle radius of RHS = 53 nm from DLS), an initial wet layerthickness of Lwet = 120 μm, and an evaporation rate E = (Lwet− Ldry)/tdry = 1.5 × 10−7 m s−1, where the wet and dry filmthicknesses are Lwet = 120 μm and Ldry = 30 μm film and thedrying time is tdry ∼ 10 min as estimated from the initial decayand/or stretched exponential fit in the drying curves reportedin Figure S5. This large Peclet number (PeNP > 1), applicablefor all the three investigated systems,23 implies a much highersolvent drying rate than particle diffusion. This imbalance intransport rates leads to an uneven packing of the latex particlesacross the film thickness during drying.21,23 Moreover, sincethe HS and S films are mostly composed by macromoleculeswith Tg lower than the film formation temperature, one has toconsider the relative contribution of particle deformation ratewith respect to solvent evaporation rate. In the framework ofRouth and Russel model,16 the parameter λ = tdef/tevap = η0RE/(γwaLwet) describing the relative interplay between solventevaporation and deformation rate of such polymer colloids canbe estimated, where η is the polymer viscosity, R the particlesize, E the evaporation rate, γwa is the water−air surfacetension, and Lwet is the initial wet layer thickness. For the Scoatings, the polymer (melt) viscosity can be assessed by usingWLF equation for acrylic coatings.68 At Τ = 23 °C, we estimateη0,S melt,dry = 1.26 × 108 Pa·s assuming a dried Tg,dry = 278 K andη0,S melt,wet ∼ 3.17 × 106 Pa·s using a wet Tg,wet = 263 K thataccounts for hydroplasticization effects69 due to partiallyretained water within the coating. Thus, λS,wet ∼ 3 × 10 −3 <λS,dry ∼ 1.1 × 10 −1 ≪ 1, suggesting that for the soft S coatingthe evaporation rate is much slower than the deformation rate.λS ≪ 1 implies wet sintering, and together with the Pe > 1, askin formation on the air−film interface of S coatings isexpected.15,16 As reported in the literature, vertical inhomoge-neity in the film structure after drying could occur.16 The largePe (>1) could result in an uneven degree of particlecoalescence, as the solvent residence time in the coatingwould differ across the coating cross section. Hence, the wet Tgof the particles (Tg,wet,S = −11.8 °C),69 in relation to theapplication (ambient) temperature, will dictate to what extentparticles will deform in the presence of water. Hence, thiscoating is expected to deform considerably but more at thebottom where water is present for a longer time. Very recently,by means of Foerster resonance energy transfer (FRET) andlight scattering, Johansmann et al. came to similar conclusionsabout a skin layer formation on the colloidal polymer filmformation.63 They proposed partial coalescence of polymerparticles into polyhedral objects.63 The same authors andothers70 also report the formation of water/surfactant arraysseparating the partially deformed and soft polymer particles intheir films. Their findings are in line with our proposition thatthe elongated nanodomains in S coatings are predominantlyoccupied by water as well as nonpolymeric cosolutes such asions and surfactants (TGA and AFM results). The scatteringpeaks observed in our GISAXS study correlate thus with the

Figure 9. Annealing effects. GISAXS patterns for HS coatings before(a) and after (b) EtOH annealing at αi = 0.2°, with intensity scalingfrom 0 to 10−3. GISAXS patterns for S coatings before (c) and after(d) EtOH annealing at αi = 0.22°, with intensity scaling from 0 to 5 ×10−4. GISAXS patterns for an aged (12 months at ambient conditionsafter film preparation) HS coating before (e) and after (f) thermalannealing (T = 150 °C, 2 h) at αi = 0.1°, with intensity scaling from 0to 5 × 103.

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center-to-center repeating distance between such heteroge-neous nanodomains along the film’s plane. For the HS system,using strain-controlled (strain = 0.1%; frequency ω = 1 Hz)extensional rheology on free-standing HS films (data notshown here), the HS polymer melt viscosity at Τ = 23 °C hasbeen measured to be η0,HS melt = 5.69 × 1010 Pa·s. We note thatthis value represents an effective viscosity of the melt from theheterogeneous HS particles. The individual viscosities of eachphase of the individual phases are not being used in thisparticular calculation. We thus find λHS ∼ 51 (>1 ≫ λS) thattogether with the Pe > 1 suggests capillary sintering (λ > 10) asthe probable formation mechanism for the HS coating.68

According to capillary sintering, the cause for deformation iscapillary pressure of the liquid between the particles whichinduces their deformation. A capillary sintering mechanism isin good agreement with the combined GISAXS, TGA, andmacroscopic staining results showing a more open porous HScoating structure at the surface compared to the S coating. Itshould also be considered that the presence of the hardundeformable domains in the HS particles can cause an arrestof the coalesced state earlier on during the film formation, ascompared to the S system. Concerning the H coating, onaccount of the fact that it constitutes only of higher (comparedto the film formation temperature) Tg phase, we attribute itsparticle deformation mechanism to dry sintering (η0,H ≫η0,HS melt at Τ = 23 °C). Indeed, a poor quality filmcharacterized by an homogeneous distribution of nanovoidsbetween particles is observed by GISAXS.

■ CONCLUSIONSWe present a detailed morphological and nanostructuralinvestigation of industrially relevant clearcoat films fromwaterborne polymer colloids for paint and protectiveapplications. The acquired results demonstrate a combinedapproach using X-ray scattering and microscopy to probe thequality of micrometer-sized films from waterborne polymercolloid formulations at the nanoscale level and across the filmthickness. The reported nanostructural heterogeneities indicatethat the particles are arrested in a partially coalesced state, wellbefore the stage of polymer chain interdiffusion betweenneighboring polymer colloids. Our findings show that filmsfrom both pure soft and soft−hard multiphase polymercolloids develop nanostructural heterogeneities distributedunevenly across the film thickness and concentrated towardthe air−film interface as a result of partial particle coalescence.The distribution of these heterogeneities implies verticalgradient in hydroplasticization effects having occurred earlierduring film formation. Variable-angle GISAXS results havebeen corroborated by cross-sectional AFM. The amount ofnanostructural heterogeneities was found to depend on themechanical properties of the colloids (i.e., chemical composi-tion), the colloidal architecture, and the film thickness. Theamount and size of the nanoscale heterogeneities (a) reflectthe degree of coalescence and (b) are larger for the multiphasecoatings with respect to the pure soft coatings. As theseheterogeneities possibly entrap surfactant molecules, residualwater molecules, and ionic species from salts, their presenceand their relative amount, expressed distinctly per colloidformulation, influence the coating performances in terms ofboth mechanical and chemical resistance against solventpenetration. The results presented here highlight thecorrelation between the nanostructure of waterborne coatingsand the macroscopic properties such as staining behavior.

Solvent and thermal annealing have been robust ways todiminish such nanostructural heterogeneities from the coatingfilms. For certain applications such as protective coatings, theseheterogeneities could be considered as defects or weak points,and thus our film post-treatment methodology provides atemplate on how to monitor and tune their contribution. Insummary, our work presents a systematic characterization ofthe waterborne coating structure at the nanoscale, relevant forthe optimization of industrially relevant coatings, including inkprimers and paints, where film thicknesses can extend toseveral micrometers. The study of these nanostructures hasbeen overlooked so far due to the lack of versatilecharacterization techniques that can operate in the nativecoating state (supported films). Moreover, our approach canbe used to inspect the nanoscale morphology at large X-raypenetration depths ξp (up to αi ∼ 4αc) of these high-qualityfilms and can also be employed in the near future to perform insitu experiments on film formation from paints, adhesives, andprotective coatings during drying.

■ ASSOCIATED CONTENT

*S Supporting InformationThe Supporting Information is available free of charge on theACS Publications website at DOI: 10.1021/acsapm.9b00601.

Differential scanning calorimetry (DSC) results withheat flow curves for the examined formulations,transmission electron microscopy (TEM) on the driedHS latex particles, SAXS modeling of suspension data,film thicknesses and dependence on particle content,kinetics of film drying, scanning electron microscopy(SEM) on dried films, Stylus profilometry, calculation ofpenetration depth for X-rays, characteristic features ofGISAXS patterns, impact of film’s edge geometry,GISAXS I(qy) cuts for the H film slightly above theYoneda position, GISAXS I(qy) cuts for the S, HS, andH films at the specular position, comparison with SAXSfrom free-standing HS and S films, GISAXS simulations,AFM after rinsing, thermogravimetry (TGA) results,GISAXS I(αf) cuts for the S, HS, and H films at αi =0.15°, GISAXS I(qy) cuts for the S, HS, and H films (i)before and after EtOH annealing without aging and (ii)before/after thermal annealing on an aged sample(PDF)

■ AUTHOR INFORMATION

Corresponding Author*E-mail g.portale@rug.nl.

ORCIDGiuseppe Portale: 0000-0002-4903-3159Author ContributionsThe project was conceived by G.P. The GISAXS measure-ments were conducted by A.V. and analyzed by A.V. and G.P.The cs-AFM experiments and the respective data analysis havebeen performed by Q.C. D.H.-M. assisted in the synchrotronexperiments. G.t.B. assisted with the SEM measurements.Synthesis and characterization of colloids were performed inthe DSM laboratories by J.S. The manuscript was written byA.V. and G.P. with the contribution of all authors. All authorshave given approval to the final version of the manuscript.

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FundingG.P. and A.V. received funding for this project by the DutchPolymer Institute (DPI) under project 914ft16.NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSThe work of A.V. and G.P. forms part of the researchprogramme of the Dutch Polymer Institute (DPI), project914ft16. We thank the DUBBLE team at the ESRF and theSpanish Beamline (ALBA Synchrotron Light Source) fortechnical support as well as Haike Ruijters (Anton Paar) for hisguidelines on zero shear viscosity measurements and SvenBroekman for support on extensional rheology measurements.Marc Stuart (RUG) is acknowledged for the TEM measure-ments. We also thank Jur van Dijken (RUG) for his support onTGA measurements. Ron Peters (DSM Coating Resins),Prof. Wesley Browne (RUG), and Simon Gree (IS2M,Mulhouse) are gratefully acknowledged for helpful discussions.

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